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Optimization of processing-microstructure-properties relationship in friction-stir welded 6061-T6 aluminum alloy
Materials Science & Engineering A 662 (2016) 136–143
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Optimization of processing-microstructure-properties relationship in friction-stir welded 6061-T6 aluminum alloy Sergey Malopheyev a,n, Igor Vysotskiy a, Vladislav Kulitskiy a, Sergey Mironov b, Rustam Kaibyshev a a b
Laboratory of Mechanical Properties of Nanostructural Materials and Superalloys, Belgorod State University, Pobeda 85, Belgorod 308015, Russia Departament of Materials Processing, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan
art ic l e i nf o
a b s t r a c t
Article history: Received 12 January 2016 Received in revised form 11 March 2016 Accepted 14 March 2016
In this work, a simple but effective approach for improvement of strength of friction-stir welded 6061-T6 aluminum alloys was elaborated. It involves friction-stir welding (FSW) at relatively high tool rotational speed and welding speed followed by standard post-weld aging. The selected combination of FSW parameters provides high welding temperature as well as rapid cooling rate. This leads to complete dissolution of strengthening precipitates in stir zone but hinders their coarsening in heat-affected zone. In turn, this enables to avoid the formation of softened region during subsequent aging and thus substantially enhances strength and improves ductility. & 2016 Elsevier B.V. All rights reserved.
Keywords: Al-Mg-Si alloys Friction-stir welding Precipitation strengthening Postweld heat treatment
1. Introduction Increasing demand for reducing weight in the automotive industry has prompted a significant interest in applications of lightweight materials, particularly aluminum alloys. Due to an attractive combination of moderate strength, excellent corrosion properties and relatively low cost, heat-treatable wrought 6xxx aluminum alloys appear to be promising candidates. However, conventional fusion welding, which is currently applied in this industry, leads to the dissolution of strengthening precipitates in these materials, thereby causing an unacceptable loss of properties. The postweld aging treatment has been demonstrated to not be effective because of the formation of solute segregations at the dendrite boundaries [1]. To overcome this problem, prior solution heat treatment is necessary but may not be practical in real production. Friction-stir welding (FSW), a solid-state process, prevents the formation of solidification structures and therefore is widely considered to be a very promising joining technique. In the 6xxx aluminum alloys, FSW also induces complex precipitation phenomena [e.g., 2–7]. Due to the large temperature gradient associated with FSW, different microstructural regions inevitably n
Corresponding author. E-mail addresses: [email protected] (S. Malopheyev), [email protected] (I. Vysotskiy), [email protected] (V. Kulitskiy), [email protected] (S. Mironov), [email protected] (R. Kaibyshev). http://dx.doi.org/10.1016/j.msea.2016.03.063 0921-5093/& 2016 Elsevier B.V. All rights reserved.
experience either precipitation dissolution or precipitation coarsening [4]. The latter effect is considered to be particularly undesirable because subsequent aging often cannot recover the material strength [4,8–11]. This phenomenon gives rise to relatively low joint efficiency even in the postweld-aged condition. To produce high-strength welds in the 6xxx aluminum alloys, it is necessary to enhance the dissolution process but minimize the coarsening process. In this case, the FSW thermal cycle would be similar to the solution heat treatment and therefore mechanical properties can be essentially restored during subsequent aging. Particle dissolution can be achieved relatively simply by increasing the welding temperature above 500–550 °C (773–823 K) by appropriately adjusting the tool rotational speed. In the 6056 aluminum alloy, for instance, this can occur at a tool rotational speed exceeding 900 rpm [12]. However, inhibiting precipitation coarsening seems to be a more challenging problem. Because it is a diffusion-controlled process, the particle coarsening kinetic is very sensitive to the duration of the weld thermal cycle. Since the heating stage during FSW is commonly known to be very short, the precipitation growth is believed to primarily occur during the cooling step. If so, the coarsening process can be minimized by accelerating the cooling speed. During normal FSW, this can be achieved by increasing the welding speed. Therefore, it appears that a combination of relatively high tool rotational speed and welding speed during FSW followed by postweld aging may provide high-strength welds in 6xxx aluminum alloys. Moreover, this approach enables to improve production productivity. This study was performed to examine this idea.
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2. Material and Experimental Procedures The material used in this investigation was a commercial AA6061 aluminum alloy supplied as 3-mm-thick sheets. This is one of the most popular 6xxx alloys used for FSW and therefore it is believed to be representative for examination of the proposed idea. The measured chemical composition of the material was (in wt%) 0.88 Mg, 0.66 Si, 0.72 Fe, 0.26 Cu, 0.12 Mn, 0.12 Cr, 0.09 Zn, and balance Al. The material was subjected to extrusion at 380 °C (653 K) with 75% reduction. To produce the peak-aged condition prior to welding, the material was T6 tempered, i.e., solutionized at 550 °C (823 K) for 1 h, water quenched and subsequently artificially aged at 160 °C (433 K) for 8 h. The T6-tempered material was friction-stir butt welded using an AccuStir 1004 FSW machine. The welding tool consisted of a shoulder having a diameter of 12.5 mm and an M5 cylindrical pin of 1.9 mm in length. To provide full-thickness joining, a doubleside FSW was applied in mutually opposite directions. To provide a relatively high welding temperature, the maximal allowable tool rotational speed of 1100 rpm was used. To examine the effect of the welding speed, welding trials were performed at tool travel speeds of 125, 380 and 760 mm/min; further increasing the weld speed led to pronounced welding defects. A tilting angle of 2.5° was employed in all cases. The principal directions of the FSW geometry were denoted as the welding direction (WD), transversal direction (TD) and normal direction (ND). To recover the mechanical properties of the welded material, the obtained joints were artificially aged at 160 °C (433 K) for 8 hours. For comparative purposes, selected welds were alternatively subjected to the T6-tempering in the manner described above. To minimize the possible effect of natural aging, the postweld heat treatments were performed within 20 min after FSW. It should be emphasized that all microstructural observations and examinations of mechanical properties in this work were conducted on the postweld heat-treated samples. The as-welded material condition is well described in the scientific literature [4– 7] and was therefore excluded from consideration in this paper. The microstructural observations were performed using optical microscopy, electron backscattered diffraction (EBSD), and transmission electron microscopy (TEM). For optical observations, the samples were prepared using conventional polishing techniques followed by final etching in Keller's reagent. A suitable surface for the EBSD and TEM were obtained using electro-polishing in a
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solution of 25% nitric acid in methanol. Particular care was taken during machining of the TEM samples from the areas of interest. The EBSD analysis was conducted using an FEI Quanta 600 fieldemission-gun scanning electron microscope (FEG-SEM) equipped with TSL OIM™ software. The TEM study was performed using a JEM-2100EX TEM operating at 200 kV. In the EBSD maps, the low-angle boundaries (2° o θ o15°) and high-angle boundaries (θ 415°) were depicted as red and black lines, respectively. In the TEM micrographs shown throughout the manuscript, the TEM specimens were typically tilted to align the incident electron beam with a o100 4zone axis, as indicated in the attached diffraction patterns. This was done to eliminate dislocation contrast and enhance the contrast associated with the second-phase particles. To view the microstructure distribution in the welds more broadly, the microhardness profiles were measured. Vickers microhardness data were obtained by applying a load of 200 g with a dwell time of 10 s in a Wolpert 402MVD microhardness tester. To examine the mechanical properties of the welds, transverse tensile tests were used. The tensile specimens were centered at the weld line and included all characteristic microstructural zones of FSW. The arrangement and dimensions of the specimens are shown in Fig. 1. The upper and lower surfaces of the specimens were mechanically polished to achieve a uniform thickness and to prevent the surface defects from affecting the tensile properties. For comparative purposes, appropriate tensile specimens were also machined from the base material. The tension tests to failure were conducted at an ambient temperature and constant crosshead velocity that corresponds to a nominal strain rate of 10 3 s 1 using an Instron 300LX universal testing machine. Two tensile specimens were tested for each material condition.
3. Results and discussion 3.1. Base material The microstructure of the base material is shown in Fig. 2. The grain structure was dominated by coarse elongated grains that contain a developed substructure (Fig. 2a). As expected, the material contained a high density of dispersoids that were evenly distributed in the grain interior (Fig. 2b). The particles often had a characteristic coffee-bean contrast indicating a coherent
Fig. 1. Weld and specimen dimensions. See Section 3 for details.
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Fig. 4. Effect of welding speed and postweld-aging on the microhardness profiles measured across the transverse cross sections of the joints. Note: The horizontal dotted line shows the microhardness in the base material. See Section 3.2 for details.
measurements showed that these particles were enriched by Cr. The additional postweld aging of the base materials did not cause significant changes either in the second-phase particles or in the mechanical properties. For simplicity, these data were not included in this paper. 3.2. Overview of the microstructure distribution in the welds
Fig. 2. Microstructure of the base material: EBSD grain-boundary map (a) and TEM image (b). In (a), low-angle boundaries and high-angle boundaries are depicted as red and black lines, respectively. The fHABs and θav in the inserts in the top right corner are the HABs fraction and the average misorientation, respectively. In (b), the insert in the bottom left corner shows the diffraction pattern. See Section 3.1 for details.
relationship with the matrix. As follows from scientific literature [13,14], those were presumably dominated by needle-shaped β” precipitates but may also include a minor fraction of lath-shaped Q’ particles. It should be noted that the metastable β” phase is well-accepted to be the main contributor to the strength of peakaged 6xxx alloys. In addition to these dispersoids, relatively coarse Al(Fe,Mn)Si intermetallics were also found (not shown). The EDS
The low magnification overviews of the transversal cross-sections of the friction-stir welds obtained at different welding speeds are shown in Fig. 3. In the figures, AS1 (AS2) and RS1 (RS2) denote advancing and retreating sides of the first (second) FSW pass, respectively. To evaluate the microstructure distribution in the postweldaged joints, microhardness profiles were measured along the dotted line in Fig. 3, and the results are summarized in Fig. 4. For clarity, the shoulder and pin dimensions are also depicted in the figure. As a first approximation, the pin diameter delineates the stir zone, whereas the shoulder diameter indicates the region most affected by the heat. The horizontal black line indicates the average microhardness in the base material. The weld produced at a relatively low tool travel speed of 125 mm/min exhibited very inhomogeneous microhardness (i.e., virtually a microstructure) distribution. In contrast to the scientific literature [e.g., 2,8,15,16] the softened region was somewhat shifted from the shoulder edge. This observation is attributable to the
Fig. 3. Low-magnification overview of transversal cross-sections of the welds produced at a welding speed of 125 (a), 380 (b) and 760 mm/min (c). AS1 (AS2) and RS1 (RS2) denote the advancing and retreating sides of the first (second) FSW pass, respectively. The white dotted lines indicate the microhardness profiles shown in Fig. 4. The white and black rectangles show the EBSD maps and TEM images given in Figs. 5 and 7, respectively. Note: All welds are shown in the postweld-aged condition.
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Fig. 6. TEM images showing the second-phase particles in the central parts of the friction-stir welds produced at a welding speed of 125 (a) and 760 mm/min (b). The inserts in the bottom right corner show the diffraction patterns. Note: All welds are shown in the postweld-aged condition. See Section 3.3 for details.
work in this area [e.g., 4]. Remarkably, all studied welds had nearly the same microhardness. Of particular interest was the observation that the stir zone microhardness was somewhat lower than the microhardness in the base material despite applying the same aging treatment (i.e., 160 °C (433 K)/8 h). This interesting effect is discussed in the next section. 3.3. Stir zone
Fig. 5. EBSD grain-boundary maps showing the grain structure in the central parts of the friction-stir welds produced at a welding speed of 125 (a), 380 (b) and 760 mm/min (c). Note: All welds are shown in the postweld-aged condition. See Section 3.3 for details.
influence of the second FSW pass (Fig. 3). The increase of the welding speed, however, promoted substantial strengthening in the softened region and thus essentially smoothed out the microhardness distribution (Fig. 4). This encouraging result obviously supports the main idea of this work. As expected, the postweld aging essentially restored the microhardness in the stir zone. This was consistent with previous
Microhardness measurements revealed two microstructural regions of particular interest: the stir zone and the softened region. The microstructures that evolved in these areas are considered in the following two sections. The EBSD maps illustrating the grain structure in the stir zones of different welds are shown in Fig. 5. A significant grain refinement is evident compared to the base material (Fig. 2a). It is also seen that increasing the welding speed enhanced the refinement effect (Fig. 5). One possible explanation for this observation was the suppression of grain growth because of the shorter cooling cycle, as discussed in Section 1. The typical TEM images of the stir zone microstructures are shown in Fig. 6. It is clear that the postweld aging promoted reprecipitation of the fine dispersoids. The precipitates had nearly the same morphology, dimensions, and even coffee-bean coherent contrast as those in the base material (Fig. 2b). Perhaps they were
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Fig. 7. TEM images showing the second-phase particles in the softened regions of the friction-stir welds produced at a welding speed of 125 (a, b), 380 (c, d) and 760 mm/ min (e, f). The inserts in the bottom right corners of (b), (d) and (f) show the diffraction patterns. Note: All welds are shown in the postweld-aged condition. See Section 3.4 for details.
also manly composed of the β” and Q’ phases. The particle reprecipitation explains the observed recovery of the microhardness in the stir zone (Fig. 4). By comparing Figs. 2b and 6, it seems, however, that the volume fraction of the precipitates in the stir zone was somewhat lower than in the base material. If so, this may be a possible reason for the incomplete restoration of strength in the stir zone, as mentioned in the previous section. Therefore, the incomplete particles precipitation requires an explanation. Despite the assumption that the welding temperature in the stir zone was high enough, the solutionized material was not quenched but cooled in air after FSW. This may lead to the formation of solute clusters (or even partial dispersoid re-precipitation) during the weld cooling cycle. If so, this should decrease the level of supersaturation, which is the driving force for the
precipitation nucleation during subsequent aging, and therefore reduce the nucleation rate [13]. This might produce the relatively low volume fraction of the precipitates in the aged stir zone. 3.4. Softened region The TEM micrographs that illustrate the precipitation patterns in the softened regions of different welds are summarized in Fig. 7. The joint produced at the lowest welding speed was characterized by the appearance of relatively coarse rod-shaped particles (Fig. 7a). Considering the specific morphology of the particles as well as their alignment with the o100 4direction, they were presumably the β’ phase. These particles are well accepted to originate from the growth of the β” precipitates. However, the
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Fig. 8. Typical deformation diagrams showing the effect of welding speed and post-weld aging on the transverse tensile properties of friction-stir joints.
fraction of the fine coherent β” dispersoids was low (Fig. 7b). As mentioned in Section 1, these observations may be interpreted in terms of particles coarsening during the weld thermal cycle and thus lowering the precipitation potential during subsequent aging. Increasing the welding speed, however, decreased the fraction of the coarse β’ phase (Fig. 7a, c, and e) but increased the properties of the fine β” precipitates (Fig. 7b, d, f). This agrees well with the smoothing of the microhardness profiles (Fig. 4) and perhaps indicates inhibition of the particle growth during welding and thus enhancement of the re-precipitation process during aging.
Fig. 9. The picture of friction-stir welded and subsequently aged specimens pulled to failure.
3.5. Tensile behavior The effect of welding speed on the tensile behavior of the aged welds is shown in Fig. 8. The deformation diagram of the base material is also included for comparative purposes. Duplicate tests showed very similar results, and thus, only one set of the diagrams is shown in the figure for each condition. The relevant mechanical properties are summarized in Table 1. Pictures of the welded specimens after failure are illustrated in Fig. 9; for clarity, the weld zone is indicated in the specimens. It is clear from Fig. 8 that the aged joints had lower strengths than the base material. However, the increase in the welding speed significantly reduced this difference. In the weld produced at 760 mm/min, the joint efficiency was 93% (Table 1). This result agrees well with microhardness measurements (Fig. 4) and is believed to be attributable to the formation of a more homogeneous particle distribution in the welded zone, as discussed in the previous section. Another important point seen in Fig. 8 is the very low ductility of the aged welds. As follows from Fig. 9, the tensile deformation was localized near the shoulder edges area, i.e., in the softened region (Fig. 4). This agrees well with the direct measurements of the strain distribution in tensioned 6xxx friction-stir welds [10,11].
Fig. 10. EBSD orientation map showing abnormal grain growth in the stir zone after postweld T6 treatment. In the map, the grains are colored according to the orientation code shown in the top right corner. Note: The microstructure shown is from a weld produced at a welding speed of 760 mm/min.
Table 1 Mechanical properties of the base material and aged welds. Material condition
Yield strength, MPa
Ultimate tensile strength, MPa
Ductility, %
Joint efficiency for yield strength, %
Failure location
Base FSW FSW FSW
290 190 260 280
350 230 290 310
11 1.6 1.4 3.3
– 66 90 93
– Softened region Softened region Softened region
material at 125 mm/min þ aging at 380 mm/min þ aging at 760 mm/min þ aging
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tensile behavior is believed to be significantly influenced by inhomogeneous microstructure distribution and thus propensity to strain localization in the softest region. Thus the observed improvement of strength and ductility of welds with increasing of the welding speed (Fig. 8) was presumably attributable to elimination of the softened region (Fig. 4) and the respective uniformity of tensile strain distribution. 3.6. Postweld T6 treatment As demonstrated above, the optimized FSW followed by the standard aging treatment process achieved high-strength welds. For comparative purposes, selected welds were subjected to alternative postweld T6 tempering. The influence of this treatment on the microstructure and properties of the joints is summarized in this section. As follows from the EBSD map in Fig. 10, the postweld T6 tempering gave rise to abnormal grain growth in the stir zone. This phenomenon is often observed in the heat-treated friction-stir welds [e.g., 5] but its origin is still unclear. In the 6xxx aluminum alloys, the abnormal grain growth has been reported by Krishnan [20]. Despite the catastrophic grain growth in the stir zone, however, the microhardness was fully recovered and uniformly distributed across the weld (Fig. 11a). This was obviously associated with the complete precipitation of the strengthening β” precipitates during the aging step of the T6 treatment. As a result, the joint efficiency for the yield strength was nearly 100% (Fig. 11b). The weld ductility also improved, but it was lower than that of the base material (Fig. 11b). As follows from Fig. 12, the limited ductility was also related to strain localization in the stir zone, most likely due to the abnormal coarse nature of the grain structure.
4. Conclusions Fig. 11. The microhardness profiles (a) and tensile-deformation diagrams (b) illustrating the effect of postweld-aging and postweld T6 treatment on the mechanical properties of the welds. Note: The mechanical properties of a weld produced at a tool rotational speed of 760 mm/min are shown.
Fig. 12. The picture of a fractured specimen of a joint subjected to the postweld T6 treatment. Note: The weld was produced at a welding speed of 760 mm/min.
It seems, therefore, that the tensile response of the welds was governed by the behavior of the softest microstructural region. The strain localization caused low global ductility (Fig. 8, Table 1). It should be noted, nevertheless, that increasing the welding speed improved the total elongation to failure (Fig. 8, Table 1). This is thought to also be associated with the microstructure homogenization within the welds. It is interesting to note that the increase of the weld strength was accompanied by the increase of the weld ductility (Fig. 8). In general, the strength and ductility of a material are well accepted to be mutually exclusive, i.e. the material strengthening usually leads to reduction of the ductility [17–19]. In the case of welded structures (similar those studied in this work), however, the
In this work, the effect of welding speed during FSW on the subsequent aging response of an AA6061-T6 aluminum alloy was studied. Welding trials were performed at tool travel speeds of 125, 380 and 760 mm/min, the obtained joints were artificially aged at 160 °C (433 K) for 8 h, and the resulting microstructure and mechanical properties were examined using TEM, microhardness measurements and transverse tensile tests. The main conclusions derived from this work are as follows. 1) Increasing the welding speed hindered precipitation coarsening in the heat-affected zone. This improved the macro-scale uniformity of the precipitation distribution in the weld zone during subsequent aging and thus enabled substantial recovery of the weld strength. Specifically, the joint produced with a tool travel speed of 760 mm/min demonstrated joint efficiency of the yield strength of 93%. Nevertheless, weld ductility during transverse tensile tests was found to be relatively low ( 3%) due to localization of the strain in the heat-affected zone. 2) The microhardness in the stir zone of the peak-aged welds was found to be slightly lower than in the base material in the T6 temper. This effect was attributed to the air-cooling conditions and the related solute segregations during the FSW cooling cycle. The segregations should reduce the precipitation potential and slow down the precipitation kinetics during postweld aging. To completely recover the material strength, longer exposure at the aging temperature is required. 3) The postweld T6 treatment causes complete recovery of the weld strength and significantly improves the weld ductility but leads to abnormal grain growth in the stir zone.
S. Malopheyev et al. / Materials Science & Engineering A 662 (2016) 136–143
Acknowledgments The financial support received from the Ministry of Education and Science, Russia, under Grant no. 14.578.21.0097 (ID number RFMEFI57814X0097) is gratefully acknowledged. The authors are grateful to the personnel of the Joint Research Centre, Belgorod State University, for their assistance with the instrumental analysis. The authors would also like to thank Marat Gazizov for assistance with the TEM observations and phase analysis.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Optimization of processing-microstructure-properties relationship in friction-stir welded 6061-T6 aluminum alloy Sergey Malopheyev a,n, Igor Vysotskiy a, Vladislav Kulitskiy a, Sergey Mironov b, Rustam Kaibyshev a a b
Laboratory of Mechanical Properties of Nanostructural Materials and Superalloys, Belgorod State University, Pobeda 85, Belgorod 308015, Russia Departament of Materials Processing, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan
art ic l e i nf o
a b s t r a c t
Article history: Received 12 January 2016 Received in revised form 11 March 2016 Accepted 14 March 2016
In this work, a simple but effective approach for improvement of strength of friction-stir welded 6061-T6 aluminum alloys was elaborated. It involves friction-stir welding (FSW) at relatively high tool rotational speed and welding speed followed by standard post-weld aging. The selected combination of FSW parameters provides high welding temperature as well as rapid cooling rate. This leads to complete dissolution of strengthening precipitates in stir zone but hinders their coarsening in heat-affected zone. In turn, this enables to avoid the formation of softened region during subsequent aging and thus substantially enhances strength and improves ductility. & 2016 Elsevier B.V. All rights reserved.
Keywords: Al-Mg-Si alloys Friction-stir welding Precipitation strengthening Postweld heat treatment
1. Introduction Increasing demand for reducing weight in the automotive industry has prompted a significant interest in applications of lightweight materials, particularly aluminum alloys. Due to an attractive combination of moderate strength, excellent corrosion properties and relatively low cost, heat-treatable wrought 6xxx aluminum alloys appear to be promising candidates. However, conventional fusion welding, which is currently applied in this industry, leads to the dissolution of strengthening precipitates in these materials, thereby causing an unacceptable loss of properties. The postweld aging treatment has been demonstrated to not be effective because of the formation of solute segregations at the dendrite boundaries [1]. To overcome this problem, prior solution heat treatment is necessary but may not be practical in real production. Friction-stir welding (FSW), a solid-state process, prevents the formation of solidification structures and therefore is widely considered to be a very promising joining technique. In the 6xxx aluminum alloys, FSW also induces complex precipitation phenomena [e.g., 2–7]. Due to the large temperature gradient associated with FSW, different microstructural regions inevitably n
Corresponding author. E-mail addresses: [email protected] (S. Malopheyev), [email protected] (I. Vysotskiy), [email protected] (V. Kulitskiy), [email protected] (S. Mironov), [email protected] (R. Kaibyshev). http://dx.doi.org/10.1016/j.msea.2016.03.063 0921-5093/& 2016 Elsevier B.V. All rights reserved.
experience either precipitation dissolution or precipitation coarsening [4]. The latter effect is considered to be particularly undesirable because subsequent aging often cannot recover the material strength [4,8–11]. This phenomenon gives rise to relatively low joint efficiency even in the postweld-aged condition. To produce high-strength welds in the 6xxx aluminum alloys, it is necessary to enhance the dissolution process but minimize the coarsening process. In this case, the FSW thermal cycle would be similar to the solution heat treatment and therefore mechanical properties can be essentially restored during subsequent aging. Particle dissolution can be achieved relatively simply by increasing the welding temperature above 500–550 °C (773–823 K) by appropriately adjusting the tool rotational speed. In the 6056 aluminum alloy, for instance, this can occur at a tool rotational speed exceeding 900 rpm [12]. However, inhibiting precipitation coarsening seems to be a more challenging problem. Because it is a diffusion-controlled process, the particle coarsening kinetic is very sensitive to the duration of the weld thermal cycle. Since the heating stage during FSW is commonly known to be very short, the precipitation growth is believed to primarily occur during the cooling step. If so, the coarsening process can be minimized by accelerating the cooling speed. During normal FSW, this can be achieved by increasing the welding speed. Therefore, it appears that a combination of relatively high tool rotational speed and welding speed during FSW followed by postweld aging may provide high-strength welds in 6xxx aluminum alloys. Moreover, this approach enables to improve production productivity. This study was performed to examine this idea.
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2. Material and Experimental Procedures The material used in this investigation was a commercial AA6061 aluminum alloy supplied as 3-mm-thick sheets. This is one of the most popular 6xxx alloys used for FSW and therefore it is believed to be representative for examination of the proposed idea. The measured chemical composition of the material was (in wt%) 0.88 Mg, 0.66 Si, 0.72 Fe, 0.26 Cu, 0.12 Mn, 0.12 Cr, 0.09 Zn, and balance Al. The material was subjected to extrusion at 380 °C (653 K) with 75% reduction. To produce the peak-aged condition prior to welding, the material was T6 tempered, i.e., solutionized at 550 °C (823 K) for 1 h, water quenched and subsequently artificially aged at 160 °C (433 K) for 8 h. The T6-tempered material was friction-stir butt welded using an AccuStir 1004 FSW machine. The welding tool consisted of a shoulder having a diameter of 12.5 mm and an M5 cylindrical pin of 1.9 mm in length. To provide full-thickness joining, a doubleside FSW was applied in mutually opposite directions. To provide a relatively high welding temperature, the maximal allowable tool rotational speed of 1100 rpm was used. To examine the effect of the welding speed, welding trials were performed at tool travel speeds of 125, 380 and 760 mm/min; further increasing the weld speed led to pronounced welding defects. A tilting angle of 2.5° was employed in all cases. The principal directions of the FSW geometry were denoted as the welding direction (WD), transversal direction (TD) and normal direction (ND). To recover the mechanical properties of the welded material, the obtained joints were artificially aged at 160 °C (433 K) for 8 hours. For comparative purposes, selected welds were alternatively subjected to the T6-tempering in the manner described above. To minimize the possible effect of natural aging, the postweld heat treatments were performed within 20 min after FSW. It should be emphasized that all microstructural observations and examinations of mechanical properties in this work were conducted on the postweld heat-treated samples. The as-welded material condition is well described in the scientific literature [4– 7] and was therefore excluded from consideration in this paper. The microstructural observations were performed using optical microscopy, electron backscattered diffraction (EBSD), and transmission electron microscopy (TEM). For optical observations, the samples were prepared using conventional polishing techniques followed by final etching in Keller's reagent. A suitable surface for the EBSD and TEM were obtained using electro-polishing in a
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solution of 25% nitric acid in methanol. Particular care was taken during machining of the TEM samples from the areas of interest. The EBSD analysis was conducted using an FEI Quanta 600 fieldemission-gun scanning electron microscope (FEG-SEM) equipped with TSL OIM™ software. The TEM study was performed using a JEM-2100EX TEM operating at 200 kV. In the EBSD maps, the low-angle boundaries (2° o θ o15°) and high-angle boundaries (θ 415°) were depicted as red and black lines, respectively. In the TEM micrographs shown throughout the manuscript, the TEM specimens were typically tilted to align the incident electron beam with a o100 4zone axis, as indicated in the attached diffraction patterns. This was done to eliminate dislocation contrast and enhance the contrast associated with the second-phase particles. To view the microstructure distribution in the welds more broadly, the microhardness profiles were measured. Vickers microhardness data were obtained by applying a load of 200 g with a dwell time of 10 s in a Wolpert 402MVD microhardness tester. To examine the mechanical properties of the welds, transverse tensile tests were used. The tensile specimens were centered at the weld line and included all characteristic microstructural zones of FSW. The arrangement and dimensions of the specimens are shown in Fig. 1. The upper and lower surfaces of the specimens were mechanically polished to achieve a uniform thickness and to prevent the surface defects from affecting the tensile properties. For comparative purposes, appropriate tensile specimens were also machined from the base material. The tension tests to failure were conducted at an ambient temperature and constant crosshead velocity that corresponds to a nominal strain rate of 10 3 s 1 using an Instron 300LX universal testing machine. Two tensile specimens were tested for each material condition.
3. Results and discussion 3.1. Base material The microstructure of the base material is shown in Fig. 2. The grain structure was dominated by coarse elongated grains that contain a developed substructure (Fig. 2a). As expected, the material contained a high density of dispersoids that were evenly distributed in the grain interior (Fig. 2b). The particles often had a characteristic coffee-bean contrast indicating a coherent
Fig. 1. Weld and specimen dimensions. See Section 3 for details.
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Fig. 4. Effect of welding speed and postweld-aging on the microhardness profiles measured across the transverse cross sections of the joints. Note: The horizontal dotted line shows the microhardness in the base material. See Section 3.2 for details.
measurements showed that these particles were enriched by Cr. The additional postweld aging of the base materials did not cause significant changes either in the second-phase particles or in the mechanical properties. For simplicity, these data were not included in this paper. 3.2. Overview of the microstructure distribution in the welds
Fig. 2. Microstructure of the base material: EBSD grain-boundary map (a) and TEM image (b). In (a), low-angle boundaries and high-angle boundaries are depicted as red and black lines, respectively. The fHABs and θav in the inserts in the top right corner are the HABs fraction and the average misorientation, respectively. In (b), the insert in the bottom left corner shows the diffraction pattern. See Section 3.1 for details.
relationship with the matrix. As follows from scientific literature [13,14], those were presumably dominated by needle-shaped β” precipitates but may also include a minor fraction of lath-shaped Q’ particles. It should be noted that the metastable β” phase is well-accepted to be the main contributor to the strength of peakaged 6xxx alloys. In addition to these dispersoids, relatively coarse Al(Fe,Mn)Si intermetallics were also found (not shown). The EDS
The low magnification overviews of the transversal cross-sections of the friction-stir welds obtained at different welding speeds are shown in Fig. 3. In the figures, AS1 (AS2) and RS1 (RS2) denote advancing and retreating sides of the first (second) FSW pass, respectively. To evaluate the microstructure distribution in the postweldaged joints, microhardness profiles were measured along the dotted line in Fig. 3, and the results are summarized in Fig. 4. For clarity, the shoulder and pin dimensions are also depicted in the figure. As a first approximation, the pin diameter delineates the stir zone, whereas the shoulder diameter indicates the region most affected by the heat. The horizontal black line indicates the average microhardness in the base material. The weld produced at a relatively low tool travel speed of 125 mm/min exhibited very inhomogeneous microhardness (i.e., virtually a microstructure) distribution. In contrast to the scientific literature [e.g., 2,8,15,16] the softened region was somewhat shifted from the shoulder edge. This observation is attributable to the
Fig. 3. Low-magnification overview of transversal cross-sections of the welds produced at a welding speed of 125 (a), 380 (b) and 760 mm/min (c). AS1 (AS2) and RS1 (RS2) denote the advancing and retreating sides of the first (second) FSW pass, respectively. The white dotted lines indicate the microhardness profiles shown in Fig. 4. The white and black rectangles show the EBSD maps and TEM images given in Figs. 5 and 7, respectively. Note: All welds are shown in the postweld-aged condition.
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Fig. 6. TEM images showing the second-phase particles in the central parts of the friction-stir welds produced at a welding speed of 125 (a) and 760 mm/min (b). The inserts in the bottom right corner show the diffraction patterns. Note: All welds are shown in the postweld-aged condition. See Section 3.3 for details.
work in this area [e.g., 4]. Remarkably, all studied welds had nearly the same microhardness. Of particular interest was the observation that the stir zone microhardness was somewhat lower than the microhardness in the base material despite applying the same aging treatment (i.e., 160 °C (433 K)/8 h). This interesting effect is discussed in the next section. 3.3. Stir zone
Fig. 5. EBSD grain-boundary maps showing the grain structure in the central parts of the friction-stir welds produced at a welding speed of 125 (a), 380 (b) and 760 mm/min (c). Note: All welds are shown in the postweld-aged condition. See Section 3.3 for details.
influence of the second FSW pass (Fig. 3). The increase of the welding speed, however, promoted substantial strengthening in the softened region and thus essentially smoothed out the microhardness distribution (Fig. 4). This encouraging result obviously supports the main idea of this work. As expected, the postweld aging essentially restored the microhardness in the stir zone. This was consistent with previous
Microhardness measurements revealed two microstructural regions of particular interest: the stir zone and the softened region. The microstructures that evolved in these areas are considered in the following two sections. The EBSD maps illustrating the grain structure in the stir zones of different welds are shown in Fig. 5. A significant grain refinement is evident compared to the base material (Fig. 2a). It is also seen that increasing the welding speed enhanced the refinement effect (Fig. 5). One possible explanation for this observation was the suppression of grain growth because of the shorter cooling cycle, as discussed in Section 1. The typical TEM images of the stir zone microstructures are shown in Fig. 6. It is clear that the postweld aging promoted reprecipitation of the fine dispersoids. The precipitates had nearly the same morphology, dimensions, and even coffee-bean coherent contrast as those in the base material (Fig. 2b). Perhaps they were
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Fig. 7. TEM images showing the second-phase particles in the softened regions of the friction-stir welds produced at a welding speed of 125 (a, b), 380 (c, d) and 760 mm/ min (e, f). The inserts in the bottom right corners of (b), (d) and (f) show the diffraction patterns. Note: All welds are shown in the postweld-aged condition. See Section 3.4 for details.
also manly composed of the β” and Q’ phases. The particle reprecipitation explains the observed recovery of the microhardness in the stir zone (Fig. 4). By comparing Figs. 2b and 6, it seems, however, that the volume fraction of the precipitates in the stir zone was somewhat lower than in the base material. If so, this may be a possible reason for the incomplete restoration of strength in the stir zone, as mentioned in the previous section. Therefore, the incomplete particles precipitation requires an explanation. Despite the assumption that the welding temperature in the stir zone was high enough, the solutionized material was not quenched but cooled in air after FSW. This may lead to the formation of solute clusters (or even partial dispersoid re-precipitation) during the weld cooling cycle. If so, this should decrease the level of supersaturation, which is the driving force for the
precipitation nucleation during subsequent aging, and therefore reduce the nucleation rate [13]. This might produce the relatively low volume fraction of the precipitates in the aged stir zone. 3.4. Softened region The TEM micrographs that illustrate the precipitation patterns in the softened regions of different welds are summarized in Fig. 7. The joint produced at the lowest welding speed was characterized by the appearance of relatively coarse rod-shaped particles (Fig. 7a). Considering the specific morphology of the particles as well as their alignment with the o100 4direction, they were presumably the β’ phase. These particles are well accepted to originate from the growth of the β” precipitates. However, the
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Fig. 8. Typical deformation diagrams showing the effect of welding speed and post-weld aging on the transverse tensile properties of friction-stir joints.
fraction of the fine coherent β” dispersoids was low (Fig. 7b). As mentioned in Section 1, these observations may be interpreted in terms of particles coarsening during the weld thermal cycle and thus lowering the precipitation potential during subsequent aging. Increasing the welding speed, however, decreased the fraction of the coarse β’ phase (Fig. 7a, c, and e) but increased the properties of the fine β” precipitates (Fig. 7b, d, f). This agrees well with the smoothing of the microhardness profiles (Fig. 4) and perhaps indicates inhibition of the particle growth during welding and thus enhancement of the re-precipitation process during aging.
Fig. 9. The picture of friction-stir welded and subsequently aged specimens pulled to failure.
3.5. Tensile behavior The effect of welding speed on the tensile behavior of the aged welds is shown in Fig. 8. The deformation diagram of the base material is also included for comparative purposes. Duplicate tests showed very similar results, and thus, only one set of the diagrams is shown in the figure for each condition. The relevant mechanical properties are summarized in Table 1. Pictures of the welded specimens after failure are illustrated in Fig. 9; for clarity, the weld zone is indicated in the specimens. It is clear from Fig. 8 that the aged joints had lower strengths than the base material. However, the increase in the welding speed significantly reduced this difference. In the weld produced at 760 mm/min, the joint efficiency was 93% (Table 1). This result agrees well with microhardness measurements (Fig. 4) and is believed to be attributable to the formation of a more homogeneous particle distribution in the welded zone, as discussed in the previous section. Another important point seen in Fig. 8 is the very low ductility of the aged welds. As follows from Fig. 9, the tensile deformation was localized near the shoulder edges area, i.e., in the softened region (Fig. 4). This agrees well with the direct measurements of the strain distribution in tensioned 6xxx friction-stir welds [10,11].
Fig. 10. EBSD orientation map showing abnormal grain growth in the stir zone after postweld T6 treatment. In the map, the grains are colored according to the orientation code shown in the top right corner. Note: The microstructure shown is from a weld produced at a welding speed of 760 mm/min.
Table 1 Mechanical properties of the base material and aged welds. Material condition
Yield strength, MPa
Ultimate tensile strength, MPa
Ductility, %
Joint efficiency for yield strength, %
Failure location
Base FSW FSW FSW
290 190 260 280
350 230 290 310
11 1.6 1.4 3.3
– 66 90 93
– Softened region Softened region Softened region
material at 125 mm/min þ aging at 380 mm/min þ aging at 760 mm/min þ aging
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tensile behavior is believed to be significantly influenced by inhomogeneous microstructure distribution and thus propensity to strain localization in the softest region. Thus the observed improvement of strength and ductility of welds with increasing of the welding speed (Fig. 8) was presumably attributable to elimination of the softened region (Fig. 4) and the respective uniformity of tensile strain distribution. 3.6. Postweld T6 treatment As demonstrated above, the optimized FSW followed by the standard aging treatment process achieved high-strength welds. For comparative purposes, selected welds were subjected to alternative postweld T6 tempering. The influence of this treatment on the microstructure and properties of the joints is summarized in this section. As follows from the EBSD map in Fig. 10, the postweld T6 tempering gave rise to abnormal grain growth in the stir zone. This phenomenon is often observed in the heat-treated friction-stir welds [e.g., 5] but its origin is still unclear. In the 6xxx aluminum alloys, the abnormal grain growth has been reported by Krishnan [20]. Despite the catastrophic grain growth in the stir zone, however, the microhardness was fully recovered and uniformly distributed across the weld (Fig. 11a). This was obviously associated with the complete precipitation of the strengthening β” precipitates during the aging step of the T6 treatment. As a result, the joint efficiency for the yield strength was nearly 100% (Fig. 11b). The weld ductility also improved, but it was lower than that of the base material (Fig. 11b). As follows from Fig. 12, the limited ductility was also related to strain localization in the stir zone, most likely due to the abnormal coarse nature of the grain structure.
4. Conclusions Fig. 11. The microhardness profiles (a) and tensile-deformation diagrams (b) illustrating the effect of postweld-aging and postweld T6 treatment on the mechanical properties of the welds. Note: The mechanical properties of a weld produced at a tool rotational speed of 760 mm/min are shown.
Fig. 12. The picture of a fractured specimen of a joint subjected to the postweld T6 treatment. Note: The weld was produced at a welding speed of 760 mm/min.
It seems, therefore, that the tensile response of the welds was governed by the behavior of the softest microstructural region. The strain localization caused low global ductility (Fig. 8, Table 1). It should be noted, nevertheless, that increasing the welding speed improved the total elongation to failure (Fig. 8, Table 1). This is thought to also be associated with the microstructure homogenization within the welds. It is interesting to note that the increase of the weld strength was accompanied by the increase of the weld ductility (Fig. 8). In general, the strength and ductility of a material are well accepted to be mutually exclusive, i.e. the material strengthening usually leads to reduction of the ductility [17–19]. In the case of welded structures (similar those studied in this work), however, the
In this work, the effect of welding speed during FSW on the subsequent aging response of an AA6061-T6 aluminum alloy was studied. Welding trials were performed at tool travel speeds of 125, 380 and 760 mm/min, the obtained joints were artificially aged at 160 °C (433 K) for 8 h, and the resulting microstructure and mechanical properties were examined using TEM, microhardness measurements and transverse tensile tests. The main conclusions derived from this work are as follows. 1) Increasing the welding speed hindered precipitation coarsening in the heat-affected zone. This improved the macro-scale uniformity of the precipitation distribution in the weld zone during subsequent aging and thus enabled substantial recovery of the weld strength. Specifically, the joint produced with a tool travel speed of 760 mm/min demonstrated joint efficiency of the yield strength of 93%. Nevertheless, weld ductility during transverse tensile tests was found to be relatively low ( 3%) due to localization of the strain in the heat-affected zone. 2) The microhardness in the stir zone of the peak-aged welds was found to be slightly lower than in the base material in the T6 temper. This effect was attributed to the air-cooling conditions and the related solute segregations during the FSW cooling cycle. The segregations should reduce the precipitation potential and slow down the precipitation kinetics during postweld aging. To completely recover the material strength, longer exposure at the aging temperature is required. 3) The postweld T6 treatment causes complete recovery of the weld strength and significantly improves the weld ductility but leads to abnormal grain growth in the stir zone.
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Acknowledgments The financial support received from the Ministry of Education and Science, Russia, under Grant no. 14.578.21.0097 (ID number RFMEFI57814X0097) is gratefully acknowledged. The authors are grateful to the personnel of the Joint Research Centre, Belgorod State University, for their assistance with the instrumental analysis. The authors would also like to thank Marat Gazizov for assistance with the TEM observations and phase analysis.
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